Handbook of Mechanical Nanostructuring

By Mahmood Aliofkhazraei

Providing in-depth information on how to obtain high-performance materials by controlling their nanostructures, this ready reference covers both the bottom-up and the top-down approaches to the synthesis and processing of nanostructured materials.
The focus is on advanced methods of mechanical nanostructuring such as severe plastic deformation, including high pressure torsion, equal channel angular processing, cyclic extrusion compression, accumulative roll bonding, and surface mechanical attrition treatment. As such, the contents are inherently application-oriented, with the methods presented able to be easily integrated into existing production processes. In addition, the structure-property relationships and ways of influencing the nanostructure in order to exhibit a desired functionality are reviewed in detail. The whole is rounded off by a look at future directions, followed by an overview of applications in various fields of structural and mechanical engineering.
With its solutions for successful processing of complex-shaped workpieces and large-scale specimens with desired properties, this is an indispensable tool for purposeful materials design.

• Language: English
• Publisher: Wiley
• Release date: May 2, 2016
• ISBN: 9783527674961

Book preview

Edited by Mahmood Aliofkhazraei

Handbook of Mechanical Nanostructuring

Volume 2

Title Page

Editor

Dr. Mahmood Aliofkhazraei

Tarbiat Modares University

Faculty of Engineering

Jalal al ahmad/Chamran highway

Faculty of Engineering

Tehran

Iran

Cover

© 1Photodiva

All books published by Wiley-VCH are carefully produced. Nevertheless, authors, editors, and publisher do not warrant the information contained in these books, including this book, to be free of errors. Readers are advised to keep in mind that statements, data, illustrations, procedural details or other items may inadvertently be inaccurate.

Library of Congress Card No.: applied for

British Library Cataloguing-in-Publication Data

A catalogue record for this book is available from the British Library.

Bibliographic information published by the Deutsche Nationalbibliothek

The Deutsche Nationalbibliothek lists this publication in the Deutsche Nationalbibliografie; detailed bibliographic data are available on the Internet at .

© 2015 Wiley-VCH Verlag GmbH & Co. KGaA,

Boschstr. 12, 69469 Weinheim, Germany

All rights reserved (including those of translation into other languages). No part of this book may be reproduced in any form – by photoprinting, microfilm, or any other means – nor transmitted or translated into a machine language without written permission from the publishers. Registered names, trademarks, etc. used in this book, even when not specifically marked as such, are not to be considered unprotected by law.

Print ISBN: 978-3-527-33506-0

ePDF ISBN: 978-3-527-67497-8

ePub ISBN: 978-3-527-67496-1

Mobi ISBN: 978-3-527-67495-4

oBook ISBN: 978-3-527-67494-7

Dedication to my Mother who teach me that the price of success is hard work

List of Contributors

Parvez Alam

Abo Akademi University

Laboratory of Paper Coating and Converting

Department of Chemical Engineering

Centre for Functional Materials

Porthaninkatu 3

20500 Turku

Finland

Igor V. Alexandrov

Ufa State Aviation Technical University

Department of Physics

12 Karl Marx Street

450000 Ufa

Russia

Mahmood Aliofkhazraei

Tarbiat Modares University

Department of Materials Science

Faculty of Engineering

14115-143 Tehran

Iran

Mohammad R. Allazadeh

Institute for Technical Physics and Materials Science

Research Centre for Natural Sciences

Hungarian Academy of Sciences

Konkoly-Thege street 29-33

1121 Budapest

Hungary

Sonia Azzaza

University 20 Août 1955 of Skikda

Department of Technology

Faculty of Technology

B.P.26 route d'El-Hadaiek

Skikda 21000

Algeria

and

Laboratory of Magnetism and Spectroscopy of Solids

Department of Physics

Faculty of Science

University Badji Mokhtar of Annaba

BP: 12

Annaba 23000

Algeria

Hamidreza Bagheri

Tarbiat Modares University

Department of Materials Science

Faculty of Engineering

14115-143 Tehran

Iran

S. Bagherifard

Politecnico di Milano

Department of Mechanical Engineering

Via La Masa 1

20156 Milano

Italy

Csaba Balazsi

Institute for Technical Physics and Materials Science

Research Centre for Natural Sciences

Hungarian Academy of Sciences

Konkoly-Thege street 29-33

1121 Budapest

Hungary

and

Institute of Materials Science and Technology

Bay Zoltán Nonprofit Ltd

for Applied Research

Fehérvári street 130

1116 Budapest

Hungary

Mariangela Bellusci

ENEA– Italian National Agency for New Technologies

Energy and Sustainable Economic Development

Technical Unit of Materials– Materials Chemistry and Technology Lab

Research Centre Casaccia

via Anguillarese 301

00123 Rome

Italy

Mohamed Bououdina

Nanotechnology Centre

University of Bahrain

Sakhir Campus

PO Box 32038

Kingdom of Bahrain

and

Department of Physics

College of Science

University of Bahrain

Sakhir Campus

PO Box 32038

Kingdom of Bahrain

Francisca G. Caballero

MATERALIA Research Group

National Center for Metallurgical Research (CENIM-CSIC)

Department of Physical Metallurgy

Avenida Gregorio del Amo, 8

28040 Madrid

Spain

Marcello Cabibbo

Università Politecnica delle Marche

Department of Industrial Engineering and Mathematics (DIISM)

Associate Professor of Metallurgy

Via Brecce Biance

60131 Ancona

Italy

Emilie Calvié

Université de Lyon

INSA-Lyon

MATEIS UMR5510

7 Avenue J. Capelle

69621 Villeurbanne Cedex

France

Pasquale Cavaliere

University of Salento

Department of Innovation Engineering

Via per Arnesano

73100 Lecce

Italy

Debasis Chaira

National Institution of Technology

Department of Metallurgical and Materials Engineering

Rourkela

Orissa 769008

India

Helen L.-W. Chan

The Hong Kong Polytechnic University

Department of Applied Physics

Hung Hom

Kowloon

Hong Kong

ZhiGang Chen

Suzhou University of Science and Technology

Physics Department

Suzhou

China

Jérôme Chevalier

Université de Lyon

INSA-Lyon

MATEIS UMR5510

7 Avenue J. Capelle

69621 Villeurbanne Cedex

France

Francesco Delogu

Università degli Studi di Cagliari

Dipartimento di Ingegneria Meccanica

Chimica, e dei Materiali

via Marengo 2

09123 Cagliari

Italy

Stefan Diebels

Universität des Saarlandes

Department of Applied Mechanics

Institut Technische Mechanik

Campus A 4.2

66123 Saarbrücken

Germany

Zaili Dong

College of Information Science and Engineering

Northeastern University

Shenyang 110819

China

Qing Bo Du

School of Materials Science and Engineering

Changzhou University

1 Gehu Road

Changzhou 213164

China

and

Jiangsu Key Laboratory of Materials Surface Science and Technology

Changzhou University

Changzhou 213164

China

Richard Dvorsky

VŠB-Technical University of Ostrava

Department of Physics

17. listopadu 15

708 33 Ostrava

Czech Republic

Nariman Enikeev

Ufa State Aviation Technical University

Institute of Physics of Advanced Materials

Ufa

Russia

and

Saint Petersburg State University

Laboratory for Mechanics of Severe Plastic Deformation

Saint Petersburg

Russia

Claude Esnouf

Université de Lyon

INSA-Lyon

MATEIS UMR5510

7 Avenue J. Capelle

69621 Villeurbanne Cedex

France

Carlos Garcia-Mateo

MATERALIA Research Group

National Center for Metallurgical Research (CENIM-CSIC)

Department of Physical Metallurgy

Avenida Gregorio del Amo, 8

28040 Madrid

Spain

R. Ghelichi

Politecnico di Milano

Department of Mechanical Engineering

Via La Masa 1

20156 Milano

Italy

Morteza Gheytani

Tarbiat Modares University

Department of Materials Science

Faculty of Engineering

14115-143 Tehran

Iran

Mario Guagliano

Politecnico di Milano

Department of Mechanical Engineering

Via La Masa 1

20156 Milano

Italy

Fatma Hadef

Université 20 Août 1955

Laboratoire de Recherche sur la Physico-Chimie des Surfaces et Interfaces LRPCSI

BP 26, Route d'El-Hadaiek

21000 Skikda

Algeria

Jing Hu

School of Materials Science and Engineering

Changzhou University

1 Gehu Road

Changzhou 213164

China

and

Jiangsu Key Laboratory of Materials Surface Science and Technology

Changzhou University

Changzhou 213164

China

Lucile Joly-Pottuz

Université de Lyon

INSA-Lyon

MATEIS UMR5510

7 Avenue J. Capelle

69621 Villeurbanne Cedex

France

Anne Jung

Universität des Saarlandes

Department of Applied Mechanics

Institut Technische Mechanik

Campus A 4.2

66123 Saarbrücken

Germany

Dr. Swapan Kumar Karak

National Institute of Technology Rourkela

Department of Metallurgical and Materials Engineering

Assistant Professor

Rourkela

Orissa 769008

India

Pavel A. Khaimovich

Department of Low Temperatures and Condensed State

National Science Center

Kharkov Institute of Physics and Technology, 1

Akademicheskaya Street

Kharkov 61108

Ukraine

Radim Kocich

Vysoká Škola Báňská – Technical University ( VŠB TU) Ostrava

Department of Materials Forming

Faculty of Metallurgy and Materials Engineering

17. listopadu 15

70833 Ostrava-Poruba

Czech Republic

and

VŠB TU Ostrava

Regional Materials Science and Technology Centre

17. listopadu 15

70833 Ostrava-Poruba

Czech Republic

Ling B. Kong

Nanyang Technological University

School of Materials Science and Engineering

50 Nanyang Avenue, 639798

Singapore

Jurij Koruza

Technische Universität Darmstadt

Institute of Material Science

Alarich-Weiss-Straße 2

64287 Darmstadt

Germany

Erhardt Lach

French-German Research Institute of Saint-Louis

5 rue du Général Cassagnou

68300 Saint-Louis

France

Hulin Li

Lanzhou University

College of Chemistry and Chemical Engineering

Lanzhou 730000

China

Pavel Lukáč

Charles University

Department of Physics of Materials

Faculty of Mathematics and Physics

Ke Karlovu 5

121 16 Praha 2

Czech Republic

Chee-Leung Mak

The Hong Kong Polytechnic University

Department of Applied Physics

Hung Hom

Kowloon

Hong Kong

Barbara Malič

Jožef Stefan Institute

Electronic Ceramics Department

Jamova cesta 39

1000 Ljubljana

Slovenia

Karine Masenelli-Varlot

Université de Lyon

INSA-Lyon

MATEIS UMR5510

7 Avenue J. Capelle

69621 Villeurbanne Cedex

France

Hamidreza Masiha

Tarbiat Modares University

Department of Materials Science

Faculty of Engineering

14115-143 Tehran

Iran

Tatjana I. Mazilova

Department of Condensed Matter

National Science Center

Kharkov Institute of Physics and Technology

Akademicheskaja, 1

Kharkov 61108

Ukraine

Sylvain Meille

Université de Lyon

INSA-Lyon

MATEIS UMR5510

7 Avenue J. Capelle

69621 Villeurbanne Cedex

France

Igor M. Mikhailovskij

Department of Condensed Matter

National Science Center

Kharkov Institute of Physics and Technology

Akademicheskaja, 1

Kharkov 61108

Ukraine

Brian S. Mitchell

Tulane University

Department of Chemical and Biomolecular Engineering

300 Lindy Boggs Building

New Orleans

Louisiana 70118-5674

USA

Bohdan N. Mordyuk

Department of Solids Acoustics

Kurdyumov Institute for Metal Physics

National Academy of Sciences of Ukraine

36 Academician Vernadsky boulevard

Kyiv 03142

Ukraine

Constance Morel

WINOA

Research and Process Development Engineering

528 avenue de Savoie, BP 3

38570 Le Cheylas

France

Maxim Murashkin

Ufa State Aviation Technical University

Institute of Physics of Advanced Materials

Karl Marx street 12

450000 Ufa

Russia

and

Saint Petersburg State University

Laboratory for Mechanics of Severe Plastic Deformation

Universitetsky prospekt 28

Old Peterhof 198504

Saint Petersburg

Russia

Fanil F. Musin

Ufa State Aviation Technical University

Department of Materials Science and Physics of Metals

12 Karl Marx Street

450000 Ufa

Russia

Sanusi K. Oladele

University of Johannesburg

The Department of Chemical Engineering

Faculty of Engineering and the Built Environment

Doornfontein

Johannesburg 2028

South Africa

Amara Otmani

Université 20 Août 1955

Laboratoire de Recherche sur la Physico-Chimie des Surfaces et Interfaces LRPCSI

BP 26, Route d'El-Hadaiek

21000 Skikda

Algeria

Franco Padella

ENEA– Italian National Agency for New Technologies

Energy and Sustainable Economic Development

Technical Unit of Materials– Materials Chemistry and Technology Lab

Research Centre Casaccia

via Anguillarese, 301

00123 Rome

Italy

I. F. Pariente

University of Oviedo

Department of Material Science and Metallurgy Engineering

Campus de Viesques

33203 GijÓn

Spain

Marzia Pentimalli

ENEA– Italian National Agency for New Technologies

Energy and Sustainable Economic Development

Technical Unit of Materials– Materials Chemistry and Technology Lab

Research Centre Casaccia

via Anguillarese, 301

00123 Rome

Italy

Prompong Pienpinijtham

Chulalongkorn University

Sensor Research Unit

Department of Chemistry

Faculty of Science

254 Phayathai Road

Pathumwan

Bangkok 10330

Thailand

Georgiy I. Prokopenko

Department of Solids Acoustics

Kurdyumov Institute for Metal Physics

National Academy of Sciences of Ukraine

36 Academician Vernadsky boulevard

03142 Kyiv

Ukraine

B. Ravisankar

National Institute of Technology

Department of Metallurgical and Materials Engineering

Tiruchirappalli 620 015

India

Julien Réthoré

Université de Lyon

INSA-Lyon

LaMCoS UMR5259

18 Rue Des Sciences

69621 Villeurbanne Cedex

France

Carlo Ricci

Università degli Studi di Cagliari

Complesso Universitario di Monserrato

Dipartimento di Fisica

S.P. Monserrato-Sestu km 0.7

09042 Cagliari

Monserrato Italy

Tadej Rojac

Jožef Stefan Institute

Electronic Ceramics Department

Jamova cesta 39

1000 Ljubljana

Slovenia

and

Centre of Excellence NAMASTE

Jamova cesta 39

1000 Ljubljana

Slovenia

Alireza Sabour Rouhaghdam

Tarbiat Modares University

Department of Materials Science

Faculty of Engineering

14115-143 Tehran

Iran

Debdas Roy

National Institute of Foundry and Forge Technology

Materials and Metallurgical Engineering Department

Hatia

Ranchi

Jharkhand 834003

India

Ilchat Sabirov

IMDEA Materials Institute

Calle Eric Kandel 2

28906 Madrid

Spain

Evgenij V. Sadanov

Department of Condensed Matter

National Science Center

Kharkov Institute of Physics and Technology

Akademicheskaja, 1

Kharkov 61108

Ukraine

Afolabi A. Samuel

University of South Africa

Department of Civil and Chemical Engineering

College of Science

Engineering and Technology

Private Mail Bag X6

Florida

Johannesburg 1710

South Africa

Pimthong Thongnopkun

Burapha University Chanthaburi Campus

Gems and Jewelry Research Unit

Faculty of Gems

57 M. 1 Chonlaprathan Road

Tambon Kamong

Amphoe Thamai

Chanthaburi 22170

Thailand

Jana Trojková

VŠB-Technical University of Ostrava

Department of Physics

17. listopadu 15

708 33 Ostrava

Czech Republic

Ruslan Valiev

Ufa State Aviation Technical University

Institute of Physics of Advanced Materials

Ufa

Russia

and

Saint Petersburg State University

Laboratory for Mechanics of Severe Plastic Deformation

Saint Petersburg

Russia

Yu Wang

The Hong Kong Polytechnic University

Department of Applied Physics

Hung Hom

Kowloon

Hong Kong

Kun X. Wei

School of Materials Science and Engineering

Changzhou University

1 Gehu Road

Changzhou 213164

China

and

Jiangsu Key Laboratory of Materials Surface Science and Technology

Changzhou University

Changzhou 213164

China

Wei Wei

School of Materials Science and Engineering

Changzhou University

1 Gehu Road

Changzhou 213164

China

and

Jiangsu Key Laboratory of Materials Surface Science and Technology

Changzhou University

Changzhou 213164

China

Yangjie Wei

College of Information Science and Engineering

Northeastern University

Shenyang 110819

China

Chengdong Wu

College of Information Science and Engineering

Northeastern University

Shenyang 110819

China

Xi Yao

Tongji University

Functional Materials Research Laboratory

Shanghai

China

Hongfang Zhang

Suzhou University of Science and Technology

Physics Department

Suzhou

China

Dandan Zhao

Lanzhou University

Key Laboratory of Nonferrous Metal Chemistry and Resources Utilization of Gansu Province

College of Chemistry and Chemical Engineering

Lanzhou 730000

China

Preface

The world of nanosize-affected materials can be classified into two main groups of nanomaterials and nanostructured materials. The term nanomaterial implies that the given material involves a nanodimensional structure and can be any form of nanoparticles (such as silicon nanoparticles), nanofibers (such as carbon nanofibers), and nanoplates (such as mica nanoplates or clay nanosheets). On the other hand, we have nanostructured materials. These materials have an apparent bulk form, but their constituent particles have nanodimensions. Nanostructured materials are prepared using either top-down or bottom-up approaches.

The severe plastic deformation (SPD) methods are among frequently used techniques in the last two decades for synthesis of nanostructured materials. By applying SPD techniques to a given metal, these processes lead to the fine microstructural changes in the materials and also lead to the change in their microstructure down to the nanometric scale. SPD is among the several techniques on the basis of top-down method for synthesis of nanostructured materials and surfaces.

Once a metal is subject to SPD at rather medium temperatures, its internal microstructure resists to the further internal deformation, so that a higher amount of stress is required for the further deformation. This state created in metals is known as work hardening or strain hardening. Indeed, work hardening results in the enhanced strength and hardness of deformed metal due to its deformation. In response to this increase in strength, metal loses its ductility and deformability. Consequently, this limitation results in the limitation of further strength rise of metal induced by applying the mechanical work, thus it results in the material failure. Many industrial deformation process conducted on metals can promote the strength of given metal due to this limitation.

The severe deformation is called to a set of methods through which it is possible to insert a relatively high mechanical work to the metal without developing any crack and failure in it. The term severe is called to this process because of the fact that the performed operations insert a severe deformation to the metal as compared to the other common metal deformation methods. The noticeable point of the SPD is that the entire process not only results in the enhanced strength of the metal, but occasionally leads to the drop in ductility and the increased softening. This difference is the exact distinction between this method and other plastic deformation methods. The governing mechanism of this property can be explained by the developed nanostructure in the metal through the SPD methods.

Considering the product geometry in this process, the SPD methods can be classified into three techniques applied on bulk, plate, and tube materials. Although these processes are not significantly different in terms of fine structural variations and all are affected by the SPD, distribution of stress and strain fields in the metal would be different due to the developed deformation in the metal. High-pressure torsion (HPT), equal channel angular pressing (ECAP), cyclic extrusion compression (CEC), accumulative roll bonding (ARB), surface mechanical attrition treatment (SMAT), and different kinds of milling are among the successfully applied SPD processes for the fabrication of nanocrystalline materials.

This handbook collects chapters about mechanical properties of nanomaterials and also important mechanical nanostructuring methods. I like to appreciate all of the contributors to this handbook and thank them for their high-quality manuscripts. I wish this collection help all researchers to benefit and develop this important type of fabrication methods of nanostructured materials.

Winter 2015

Mahmood Aliofkhazraei

Part I

Mechanical Properties of Nanostructured Materials

1

Mechanical Properties of Nanocrystalline Materials

Pasquale Cavaliere

1.1 Introduction

Nanostructured materials attracted a wide scientific interest in the past decade. The strength of metals and alloys is strongly influenced by the grain size. The attractive properties of nanocrystalline (NC) metals and alloys are the high yield and fracture strength, the improved wear resistance, and the superplastic behavior at relatively low temperatures and high strain rates as compared to microcrystalline (MC) materials. NC metals also exhibit high strain rate sensitivity as compared to MC materials [1, 2]. The strength of the metals is related to the microstructure as described by the well-known Hall–Petch (H-P) relationship. Generally, it is observed that the rate of strength increases by decreasing the mean grain size below 100 nm and the strength decreases by decreasing the grain size below about 20–10 nm mean grain size; such a behavior has been commonly indicated as H-P breakdown, implying a transition in the deformation modes of metals by decreasing the grain size from NC range down to very low levels. Recent investigation has suggested that dislocation-accommodated boundary sliding is the main deformation process governing the entire deformation in NC metals [3]. Actually, different processing methods are available to produce ultrafine-grained materials (UFG), such as mechanical alloying (at room and low temperatures) with consequent consolidation (compaction and/or extrusion) and severe plastic deformation (SPD) (high-pressure torsion [4], HPT or equal-channel angular pressing [5], ECAP), generally leading to the production of UFG materials [6, 7] and gas-phase condensation of particles with consequent consolidation or electrodeposition capable of producing metals in the range of NC grain size. SPD is useful in producing bulk materials with enhanced strength, hardness, and wear and superplastic properties at relatively a low temperature and high strain rates. Some of the general properties in the available literature are summarized in Table 1.1. The mechanisms of deformation and the properties of the material not only depend on the average grain size but are also strongly influenced by the grain size distribution and grain boundary structure (e.g., low-angle versus high-angle grain boundaries). The wide application of UFC NC metals in the modern industry is related to the increased understanding of their damage resistance and of the mechanical mechanisms involved in the deformation, particularly under cyclic loading. As a general behavior, it was observed that the fatigue limit of NC metals increases with decreasing grain size, and the crack initiation susceptibility decreases with increasing crack growth rate coupled with grain refinement [8]. The main damage mechanism has been recognized in the early strain localization and microcrack formation for ECAP materials. In general, a high decrease in the fatigue properties was shown for SPD materials in the low cycle fatigue (LCF) regime of intermediate-to-high plastic strain amplitudes. On the contrary, in the high cycle fatigue (HCF) regime of intermediate-to-low plastic strain amplitudes, it results in high enhancement of the fatigue resistance for materials with grain refinement. In MC materials, a reduction in the grain size generally results in an increase in strength, which engenders an increase in the fatigue endurance limit during stress-controlled cyclic loading of initially smooth-surfaced laboratory specimens.

Table 1.1 Properties of ultrafine and NC materials produced via different techniques

As the total fatigue life under the aforementioned conditions is dominated by crack nucleation and as the fatigue cracks generally nucleate at the free surface, grain refinement is considered to result in improvements in fatigue life and endurance limit, with all other structural factors set aside. On the other hand, a coarse grain structure with lower strength and enhanced ductility generally plays a more beneficial role in the strain-controlled fatigue response of metals and alloys. It should be noted, however, that it is often difficult to isolate the sole effects of grain size on fatigue response because other structural factors such as precipitate content, size and spatial distribution, stacking fault energy and the attendant equilibrium spacing of partial dislocations, and crystallographic texture are also known to have an important effect on the fatigue characteristics of MC metals. In NC materials with finest grains, plastic flow is conducted mostly by grain boundary processes. In NC materials with intermediate grains, plastic flow is often conducted by both lattice dislocation slip and grain boundary processes. However, if plastic flow and diffusion are not intensive in NC materials with intermediate grains, and/or these materials contain pre-existent nanocracks and pores, brittle fracture tends to occur. Dimpled rupture, dislocation activity at the crack tip, and formation of voids at grain boundaries and triple junctions in the regions ahead of the advancing crack were observed. In the early stages of deformation, dislocations are emitted from the grain boundaries under the influence of the applied stress. Triple junction voids and wedge cracks can also result from grain boundary sliding if the resulting displacements at the boundary are not accommodated by diffusional or power-law creep. These grain boundary and triple junction voids then act as sites for nucleation of the dimples. The deformation and fracture processes are closely related to the coupling of dislocation-mediated plasticity and formation and growth of voids. Irrespective of the fracture mechanism, it is evident that the fracture is heavily influenced by the microstructural features such as the presence of nanoscale voids or bubbles and the presence of grown-in twins, which have, so far, been mostly neglected [9]. The presence of grown-in twins has been suggested as an interface control mechanism in coarse-grained metals, and they represent a relevant microstructural feature that influences fracture, as many of the NC metals contain grown-in twins. The aim of this chapter is to provide a deep and complete understanding on the overall microstructural and mechanical properties of nanostructured metals and alloys.

1.2 Static Properties

1.2.1 Tensile Behavior

The strength of metals is related to the microstructure as described by the well-known H-P relationship. Generally, it is observed that the rate of strength increases by decreasing the mean grain size below 100 nm, and the strength decreases by decreasing the grain size below about 20–10 nm mean grain size; such a behavior has been commonly indicated as H-P breakdown, implying a transition in the deformation modes of metals by decreasing the grain size from NC range down to very low levels. In these alloys produced via SPD, which leads to materials characterized by ultrafine grains, the yield and ultimate tensile strengths increase with decreasing grain size; such an increase in yield is generally related to the deformation level such as the number of passes in ECAP (Figure 1.1), and such an increase in strength is generally coupled with a decrease in the material ductility [10].

nfg001

Figure 1.1 Stress versus strain at room temperature when testing under tension at an initial strain rate of 3 × 10−3 s−1: curves are shown for the as-received AA5754 and after pressing through one to six passes.

As clearly shown in Figure 1.2, aluminum alloys subjected to SPD via ECAP exhibit a different strain softening as a function of the number of passes. The stability of the mechanical properties is governed by precipitation, and in the alloys strengthened by stable precipitates inhibiting grain growth by impeding grain boundary mobility, the materials exhibit finer structures at the same level of SPD. An interesting example is shown for AA6XXX reinforced with Sc–Zr and produced via ECAP (Figure 1.2). After SPD and aging, the 6106 Sc alloy is stronger than the 6106 Zr alloy under the same condition, because of the more effective precipitation hardening in the Sc–Zr-modified alloy [11]. The Al3(Zr1−x,Scx) precipitates have been demonstrated to provide more effective hardening and are more stable at high temperatures as compared to Al3Zr. In addition, the effectiveness of the different particles on grain refinement is stronger in the case of Sc-modified alloys (170 µm after four passes) as compared to the Zr-modified alloy (200 µm after four passes).

nfg002

Figure 1.2 Tensile response of the 6106 Zr and 6106 Sc alloys after solution treatment + ageing, and solution treatment + ECAP + aging conditions.

As in the case of 5754 alloy, after ECAP, the alloys exhibit strain softening leading to failure after reaching the maximum tensile peak. Such a behavior is normally observed in very complex alloys, whereas it is not observed in pure metals such as Ni produced via electrodeposition (Figure 1.3a), and it can also be explained in terms of dislocation generation and rearrangement. It is clear that this phenomenon is much more pronounced in the materials obtained through SPD. In fact, also in the case of Ni–W alloys produced via electrodeposition (Figure 1.3b), we can underline the absence of softening with strain, demonstrating that such a behavior is related to the dislocation density and energy being much higher than those in the materials produced via SPD.

nfg003

Figure 1.3 Tensile behavior of pure Ni with a different grain size produced via electrodeposition (a), Ni–W alloy with a different grain size (b).

It should be underlined that the very low level of macroscopic tensile ductility in NC fcc metals is due to the localization of deformation demonstrated by the observation of the fracture surface in the tensile tested specimens. In such materials, in fact, the fracture surfaces comprise dimples larger than the original grain size, and the number of dimples increases with decreasing material grain size. In general, while electrodeposited NC Ni and Ni–W alloys exhibit deformation behavior with decrease in grain refinement down to a quasi-absence due to reduction in the fractured area, in UFC materials, a detectable necking is observed. For electrodeposited NC metals, the strength increases with decreasing grain size. Strain softening is not observed in these metals, owing to the lower energy of dislocations with respect to severely deformed materials such as those produced via ECAP. Recovery is not observed in deformed NC metals. On the contrary, room temperature recovery can be demonstrated in ultrafine materials produced via ECAP or HPT. The deformation mechanism, ductility, hardening–softening behavior, and strain rate sensitivity are strongly related to one another. These features are discussed in the next section.

1.2.2 Nanoindentation

With the development of nanostructured metals and alloys, instrumented nanoindentation seems to be very useful in obtaining the fundamental mechanical properties and for understanding the fundamental material physics. This characterization technique is a very powerful tool, because of the fact that the tested volume of material is compatible with the microstructure. Many papers were presented in the literature on nanostructured material characterization through instrumented nanoindentation; in addition to hardness and yield strength (Figure 1.4), such a characterization technique seems to be very useful in the analyses of material hardening and/or softening [12].

nfg004

Figure 1.4 Hardness and yield strength measure obtained by instrumented nanoindentation for different nanostructured materials.

This technique has also been employed to provide experimental evidence of the dynamic properties of NC materials. Nanoindentation fatigue experiments can provide very useful information on plastic zone propagation, cyclic hardening, and crack nucleation and growth in nanostructured materials. The material behavior can be explained similarly to crack propagation. In static loading, the plasticity surrounding the crack tip either blunts the crack or shields the crack tip from the external stress. Dynamic loading leads to a dynamic process between the effective applied stress and the internal stress, which is similar to the dislocation generation and annihilation to the crack tip in the conventional crack propagation tests [13]. Multistep nanoindentation is an interesting technique. A fixed strain (in terms of indentation depth) is reached in a single indentation or through increasing levels of deformation, indicating a variation in the mechanical properties of the material by a difference in the hardness values implying the hardening or softening behavior of the tested material (Figure 1.5). Another fundamental aspect is represented by the possibility of calculating the material strain rate sensitivity through the loading rate variation during nanoindentation. High strain rate sensitivity can lead to a general improvement in the strength and ductility properties of the materials. A deeper understanding of this aspect can provide important information on the mechanical evolution of a large variety of NC metals for engineering applications.

nfg005

Figure 1.5 Nanoindentation behavior for pure electrodeposited Ni showing hardening (a), for pure electrodeposited Co showing softening (b).

The strain rate sensitivity of a material is defined as the variation in the flow stress with the strain rate at a given level of strain for a fixed temperature, and it can be expressed as:

equation

where k is the Boltzmann constant, T is the absolute temperature, σ is the flow stress, and v* is the activation volume, which can be considered as the derivative of the activation energy with respect to the effective shear stress. By employing nanoindentation measurements, the flow stress can be related to the measured hardness (H = 3σ). The calculation of strain rate sensitivity is crucial in revealing many deformation mechanisms in nanostructured metals and alloys (Figure 1.6). Here, it is evident that a reduction in grain size from micro- to NC regime leads to an increase of an order of magnitude of strain rate sensitivity of plastic deformation. In general, the different behaviors observed for UFG and NC metals can be explained in terms of activation volume. A small activation volume of dislocation mobility is responsible for the variation in strain rate sensitivity with decreasing mean grain size of the metals. In addition, the mechanisms of dislocation generation at the grain boundaries coupled with grain rotation and migration are responsible for the entire plastic deformation in NC metals. Such mechanisms disappear by increasing the grain size from NC to UFG regime, thus decreasing the strain rate sensitivity of the materials at room temperature.

nfg006

Figure 1.6 Room-temperature strain rate sensitivity of electrodeposited Ni with a different grain size.

1.3 Wear Properties

As in the case of tensile properties, wear behavior in metals is influenced by grain size [14]. The normal indentation test represents a limited application in predicting the tribological response. On the other hand, the scratch test, in which a hard indenter is slid across the surface of the material, is a tool for testing materials under conditions of controlled abrasive wear [15]. Frictional sliding or scratch testing is an alternative technique to characterize the hardness and response of materials in contact with hard indenters. During frictional sliding under conditions where the tip apex angle is sufficiently large to prevent the onset of discontinuous plasticity, a steady-state regime is reached after applying a constant normal force over a sufficient distance.

The characteristics of the residual scratch profile in this steady-state regime can be used to document the resistance and properties of the materials. From the scratch tests, it can be underlined that, with decreasing grain size, the material shows a significant increase in the friction coefficient (Table 1.2). Such a behavior becomes more evident by increasing the applied load and the penetration depth due to the different strain hardening of the material, which increases with decreasing mean grain size. In scratch tests, the normalized hardness and the normalized pile-up height are sufficient to determine the plastic strain hardening exponent and the initial yield strength. The friction coefficient decreases with increasing yield strength and strain hardening; however, the yield strength, due to grain refinement, appears more effective in improving the wear properties of the materials.

Table 1.2 Variation of friction coefficient with scratch load and penetration depth for electrodeposited nickel

1.4 Fatigue Properties

The fatigue properties of materials are strongly governed by the grain size variation. Many experimental evidences can be presented both in the ultrafine and in the NC regime [16–18]. The first evidence can be underlined from the S–N curves of materials (Figure 1.7). In general, grain refinement via ECAP leads to an increase in fatigue properties in stress-controlled tests. The main damage mechanism has been recognized in the early strain localization and microcrack formation for the ECAP materials. In general, a strong decrease in fatigue properties was shown for SPD materials in the LCF regime of intermediate-to-high plastic strain amplitudes. On the contrary, in the HCF regime of intermediate-to-low plastic strain amplitudes, it results in high enhancement of the fatigue resistance for materials with grain refinement. In addition, it was observed that an annealing treatment, subsequent to the ECAP process, lead to enhancement of the LCF properties due to the increase in ductility. Such a behavior is obtained by partially recovering the grain boundary region that has been heavily distorted during processing. Pure UFG Ti, Al, and Ni show a decrease in ΔK threshold and an increase in crack propagation rate. On the contrary, pure Cu revealed higher susceptibility to crack initiation and a faster crack growth rate. Such a behavior is due to the different crack path related to the ductility variation after SPD demonstrated by the strain-controlled fatigue tests. For electrodeposited NC and UFG Ni, it was observed that the fatigue behavior of the materials is highly strain dependent. Even if grain refinement leads to an increase in the number of cycles to failure at the same stress levels investigated, the results for very close microstructures (20 and 40 nm) resulted in a strong function of the ductility (very high stresses).

nfg007

Figure 1.7 Fatigue curves of Ti produced via severe plastic deformation with UFC microstructure (a) and of Ni with ultrafine and NC microstructures (b).

In addition, by analyzing the strain amplitude as a function of the number of cycles to failure for all the electrodeposited materials, it can be concluded that the sensitivity to cyclic hardening increases by decreasing the material mean grain size.

1.5 Crack Behavior

Grain refinement, due to SPD, produces a decrease in ΔK threshold and an increase in crack propagation rate. The primary mechanism responsible for the accelerated fatigue crack growth rate observed with decreasing grain size is the reduction in crack path deflection with grain refinement. Microstructural size scales can play a dominant role in crack morphology and in the fracture mode, particularly near the threshold regime. Periodic deflections in the fatigue crack at the grain boundaries during crystallographic fracture can lead to a relatively tortuous crack path in coarser-grain materials. ECAP Ti exhibits a straight crack path, as compared to its MC counterpart (Figure 1.8a). The fatigue crack rate as a function of ΔK for all the Ni materials at a load ratio of R = 0.25 is shown in Figure 1.8b. At all the investigated stress levels, the material is less sensitive to crack initiation with decreasing mean grain size. On the other hand, the resistance to crack growth decreases with grain refinement. The fatigue crack rate as a function of ΔK for Co-based materials at a load ratio of R = 0.25 is shown in Figure 1.8c.

nfg008

Figure 1.8 Crack growth behavior as a function of ΔK for the ECAP and MC pure Ti (a); crack growth rate as a function of ΔK for the electrodeposited pure Ni in the NC and UFG regimes (b); for NC Co as compared with its MC counterpart (c).

As a general trend, the crack growth rate is governed by the crack path. In the NC Co-based materials, the path appears very flat, and it is governed by the local brittleness of such NC metals, while in the case of MC materials, the path appears completely different with localized ductile. It was described that the possible intergranular fatigue crack growth behavior could be due to the high dislocation density in UFG materials, coupled with the presence of nonequilibrium grain boundaries due to trapping and accommodation of lattice dislocations during SPD. Small-sized grains lead to reduced flaw sizes and increased difficulties for the imposed stress concentration at the flaw to exceed the critical toughness of the material, thus suppressing early crack nucleation and propagation. In the grain size range of 100–500 nm, the deformation mechanisms are similar to those in fine-grained traditional materials; whereas for grain sizes in the range of 50–100 nm, dislocations are emitted from, and are annihilated at, the grain boundaries; in the grain size range of 10–50 nm, partial dislocation emission and deformation twinning constitute the major deformation mechanisms; and for grain sizes below 10 nm, grain boundary sliding is the dominant deformation mechanism [19]. By analyzing the behavior of crack tip plasticity on the change in the near-tip field quantities of the plastic zone in different NC alloys with different configurations of grain size distribution, it was possible to study the problem of a crack approaching the interfaces at which the plastic properties of the material change continuously and linearly [20]. It was observed that the J-integral decreases when passing from the harder to the softer material, demonstrating that the potential energy for crack propagation increases in the negative plastically graded configuration (in which the grain size linearly varies between 20 and 100 nm from the surface to the bottom ), while the potential energy decreases in the case of positive graded configuration (in which the grain size linearly varies between 100 and 20 nm from the surface to the bottom). The J-integral variation in the negative and positive configurations, as compared with the constant 100 nm grain size for Ni–W electrodeposited alloy, is shown in Figure 1.9. From this study, it can be concluded that the graded properties of the electrodeposited alloys play a very important role in the control of crack growth, and particularly, the negative graded configuration appears very useful in reducing the crack growth rate.

nfgz009

Figure 1.9 Comparison between the calculations of the J-integral for the negative and positive plastically graded sheets.

From the engineering point of view, the aforementioned result is very useful; in fact, it is demonstrated that, for NC metals, the susceptibility to crack initiation decreases with decreasing grain size while it increases the crack propagation rate. By producing structures with negative graded configuration, it is possible to obtain a surface with very low susceptibility to crack initiation and a bulk structure in which the crack growth rate is continuously decreased in each section.

1.6 Conclusions

The analyses of the microstructural and mechanical properties of nanostructured materials lead to very bright horizons to new researches and industrial applications. This chapter describes the increase in the mechanical properties of nanostructured metals and alloys by increasing the grain refinement up to a level at which the grain size volume begins influencing the dislocation generation and motion, leading to an inversion in such a behavior, well known as Hall–Petch inversion. The microstructural and mechanical features of nanostructured materials strongly depend on the production techniques. Actually, UFG metals produced via SPD exhibit a refining limit of few hundreds of nanometers depending on the total deformation strain. Such severe deformation leads to a microstructure characterized by a high dislocation energy level, which leads to a material very sensitive to room- and high-temperature deformation and to the microstructure modifications as a consequence of heat treatments. Grain refinement can be improved by employing different processing techniques such as electrodeposition. Such techniques can obtain pure bulk metals characterized by grain sizes below 10 nm without defects. A decrease n the grain size down to such levels has a strong effect on the increase in strength but coupled with a large reduction in ductility. Additionally, the hardening behavior is strongly influenced by the grain size at such refining levels. These aspects are well evidenced by employing a characterization technique such as instrumented nanoindentation, which is capable of probing nanostructured materials in a broad range of forces, strains, and strain rates. This technique was revealed to be very useful in measuring the strain rate sensitivity of nanostructured metals in a broad range of grain sizes, leading to the definition of deformation mechanisms during strain in NC metals and alloys. Very useful information is obtained from the study of fatigue properties of aforementioned class of new materials. UFG materials produced via SPD show an increase in fatigue limit with decreasing grain size in the high cycle regime; the low cycle behavior is strongly influenced by dislocation mechanisms such as large deformation recovery. For the materials produced via SPD, even if grain refinement leads to an increase in the number of cycles to failure at the same stress levels investigated, the results for very close microstructures (20 and 40 nm) resulted in a strong function of the ductility (very high stresses). In addition, by analyzing the strain amplitude as a function of the number of cycles to failure for all the electrodeposited materials, it can be concluded that the sensitivity to cyclic hardening increases by decreasing the material mean grain size. For such materials, produced via electrodeposition, the fatigue crack growth tests, performed over a broad range of stress levels, revealed that these materials are less sensitive to crack initiation with decreasing mean grain size while the resistance to crack growth decreases with grain refinement. A very interesting perspective is achieved by the possibility of producing plastically graded bulk structures via electrodeposition. By tuning the electrodeposition current and the bath temperature, it is possible to vary the material composition in terms of the alloying elements, consequently varying the grain size along with the thickness of the thin films. The analyses of fatigue properties of such structures showed the possibility to control the grain size initiation and growth through the control of grain size and distribution along the crack paths.

References

1. Gleiter, H. (2000) Nanostructured materials: basic concepts and microstructure. Acta Mater., 48, 1–29.

2. Meyers, M.A., Mishra, A., and Benson, D.J. (2006) Mechanical properties of nanocrystalline materials. Prog. Mater Sci., 51, 427–556.

3. Cavaliere, P. (2008) Strain rate sensitivity of ultra-fine and nanocrystaline metals and alloys. Physica B, 403, 569–575.

4. Valiev, R.Z., Islamgaliev, R.K., and Alexandrov, I.V. (2000) Bulk nanostructured materials from severe plastic deformation. Prog. Mater Sci., 45, 103–189.

5. Valiev, R.Z. and Langdon, T.G. (2008) Using high-pressure torsion for metal processing: fundamentals and applications. Prog. Mater Sci., 53, 893–979.

6. Zhu, Y.T. and Lowe, T.C. (2000) Observations and issues on mechanisms of grain refinement during ECAP process. Mater. Sci. Eng. A, A291, 46–53.

7. Zhilyaev, A.P., Kim, B.-K., Szpunar, J.A., Bar'o, M.D., and Langdon, T.G. (2005) The microstructural characteristics of ultrafine-grained nickel. Mater. Sci. Eng. A, A391, 377–389.

8. Cavaliere, P. (2009) Fatigue properties and crack behavior of ultra-fine and nanocrystalline pure metals. Int. J. Fatigue, 31, 1476–1489.

9. Zhu, Y.T., Liao, X.Z., and Wu, X.L. (2012) Deformation twinning in nanocrystalline materials. Prog. Mater Sci., 57, 1–62.

10. Valiev, R.Z. and Langdon, T.G. (2006) Principles of equal-channel angular pressing as a processing tool for grain refinement. Prog. Mater Sci., 51, 881–981.

11. Sakaia, G., Horitaa, Z., and Langdon, T.G. (2005) Grain refinement and superplasticity in an aluminum alloy processed by high-pressure torsion. Mater. Sci. Eng. A, A393, 344–351.

12. Cavaliere, P. (2009) Mechanical properties of nanocrystalline metals and alloys studied via multi-step nanoindentation and finite element calculations. Mater. Sci. Eng. A, A512, 1–9.

13. Cavaliere, P. (2010) Cyclic deformation of ultra-fine and nanocrystalline metals through nanoindentation: similarities with crack propagation. Procedia Eng., 2, 213–222.

14. Wang, L., Gao, Y., Xu, T., and Xue, Q. (2006) A comparative study on the tribological behavior of nanocrystalline nickel and cobalt coatings correlated with grain size and phase structure. Mater. Chem. Phys., 99, 96–103.

15. Cavaliere, P. and Prete, P. (2010) Tribomechanisms of pure electrodeposited Ni at ultra-fine and nanoscale level. Wear, 268, 1490–1503.

16. Mughrabi, H. and Höppel, H.W. (2010) Cyclic deformation and fatigue properties of very fine-grained metals and alloys. Int. J. Fatigue, 32, 1413–1427.

17. Vinogradov, A. and Hashimoto, S. (2001) Multiscale phenomena in fatigue of ultra-fine grained materials-an overview. Mater. Trans., 42 (1), 74–84.

18. Mughrabi, H., Höppel, H.W., and Kautz, M. (2004) Fatigue and microstructure of ultrafine-grained metals produced by severe plastic deformation. Scr. Mater., 51, 807–812.

19. Farkas, D., Willemann, M., and Hyde, B. (2005) Atomistic mechanisms of fatigue in nanocrystalline metals. Phys. Rev. Lett., 94 (16), Art. No. 165502.

20. Cavaliere, P. (2008) Crack tip plasticity in plastically graded Ni–W electrodeposited nanocrystalline alloys. Comput. Mater. Sci., 41, 440–449.

2

Superior Mechanical Properties of Nanostructured Light Metallic Materials and Their Innovation Potential

Maxim Murashkin, Ilchat Sabirov, Nariman Enikeev and Ruslan Valiev

2.1 Introduction

The name light metals has been traditionally given to Al, Ti, and Mg because they are frequently used to reduce the weight of the components and structures used in various industries such as automotive, aerospace, electrical, and structural engineering. [1]. These three metals exhibit relative densities in the range from 1.7 (for Mg) to 4.5 g cm−3 (for Ti) as compared with 7.9 g cm−3 density of the classical structural metal such as Fe [1]. In addition to sufficient mechanical strength, these light metals exhibit good functional properties. The progress in their practical application has been determined by intensive R&D works on these light metals and their alloys. The light metals display very good functional properties such as high corrosion resistance (Al and Ti) [2], good weldability (Al and Ti) [2], and good biocompatibility (Ti and Mg) [2], although their low mechanical strength was the main obstacle for their wide industrial application. In the past, alloying, precipitation hardening, and work hardening were considered as the main strategies to improve the strength of the light metals [3]. Within the past two decades, it was demonstrated that grain refinement down to nanoscale can dramatically improve their properties [4–6]. Various processing techniques were developed for the fabrication of bulk nanostructured (NS) light metals. Analysis of the main techniques is performed in Table 2.1. These processing techniques can be divided into two major groups: bottom-up and top-down approaches. Powder milling followed by consolidation of the powders is an example of the bottom-up approach[8, 9]. Greater attention was attracted by the top-down approach, which includes severe plastic deformation (SPD) techniques [10–12]. A significant advantage of the SPD techniques is their applicability to most of the light metals and their alloys [12, 13]. Nowadays, from a practical point of view, specifically SPD of bulk billets appears to be the most promising processing method for manufacturing bulk NS light metals, which are the focus of this overview.

Table 2.1 Analysis of the main methods developed for the fabrication of NS light metals

There has been a significant body of experimental and theoretical research on the mechanical properties and mechanical behavior of NS light metals. It was demonstrated that the strength, ductility, and fatigue life of these materials are primarily controlled by their microstructure, where the grain size, morphology and size of second-phase precipitates, segregations of solute atoms, and character of the grain boundaries are very important parameters. The intelligent microstructural design of the NS light metals allows for the fabrication of light metals and alloys with superior mechanical properties, which could not be imagined two decades ago. The major objective of this chapter is to overview the recent advances in the area of SPD-processed NS light metals and alloys and to highlight the modern trends in research. Emphasis is laid on the main objective to dramatically enhance the mechanical properties of the light metals, such as strength, ductility, and fatigue life via nanostructuring and on the microstructure–properties relationship. Specific attention is paid to the current industrial applications of the NS light metals and their innovation potential.

2.2 Nanostructuring of Light Metallic Materials Using SPD Methods

The SPD techniques developed for grain refinement in metals have attracted great attention due to their ability to produce significant grain refinement in bulk metallic samples, which can be further used for various engineering applications. Depending on the SPD technique and processing parameters, the microstructure can be refined down to ultrafine (0.1–1.0 µm) or nano- (<100 nm) scales. The most popular SPD methods for processing bulk NS metallic materials include equal-channel angular pressing (ECAP) and its modifications [4], high-pressure torsion (HPT) [5], accumulative roll bonding (ARB) [14, 15], friction stir processing (FSP) [16], hydrostatic extrusion (HE) [17, 18], cryorolling (CR) [19], and so on (Table 2.1). More detailed information on the existing SPD methods can be found in the recent comprehensive review [12]. In this chapter, we do not consider in detail these SPD methods, as they have already been well described in the original papers and reviews cited earlier and references therein.

Figure 2.1 illustrates some examples of ultra-fine grained (UFG) and NS microstructures developed in certain light metals via SPD. It is seen that a very homogeneous nanostructure with an average grain size of 22 nm can be developed in a highly alloyed AA7075 via HPT (Figure 2.1a) [20]. Homogeneous UFG microstructures with equiaxed grains having an average size of 200 nm can be also obtained in pure Ti if high strains are induced into the material via the combination of ECAP-Conform and drawing (Figure 2.1b) [21]. It is very difficult to refine the microstructure of Mg and its alloys because they are prone to cracking during SPD. Thus, the development of SPD techniques for Mg alloys can require significant efforts to determine the correct processing parameters [22]. For example, a multitemperature ECAP technique was developed for the AZ31 Mg alloy after many experimental trials in [22]. It consisted of four sequential ECAP steps: (1) processing at 200 °C for four passes with 10 cm min−1 pressing speed, (2) processing at 150 °C for four passes with 3 cm min−1 pressing speed, (3) processing at 125 °C for two passes with 1 cm min−1 pressing speed, and (4) processing at 115 °C for two passes with 1 cm min−1 pressing speed. The billet was rotated about the extrusion axis at 180 between steps (1), (2), and (3) and rotated at 90 between steps (3) and (4). This processing technique led to the formation of a homogeneous UFG microstructure with an average grain size of 370 nm (Figure 2.1c), exhibiting excellent mechanical properties (see Section 2.4). A similar processing technique with varying ECAP temperature, pressing speed, and rotation angle was developed for AZ31B alloy in [23] and for pure Mg in [24].

nfg001

Figure 2.1 Some examples of UFG and NS microstructures developed in light metals via SPD. (a) AA7075 after HPT for 5 turns at room temperature; (b) pure Ti (Grade 2) after ECAP-C followed by drawing; (c) Mg alloy AZ31 after multitemperature ECAP.

The transmission electron microscopy (TEM) images are reproduced from Refs [20–22] with permission of publishers.

2.3 Superior Mechanical Strength of NS Light Metals and Alloys

Five main strategies to increase the strength of light metals can be outlined.

Solid solution strengthening. The strengthening effect in the presence of solute atoms in the matrix depends mainly on the concentration of solute atoms dissolved in the matrix and the difference between the radius of the solute atoms and that of the matrix atoms [25]. Recent investigations on the Al alloys using the atom probe tomography (APT) technique have shown that the solute atoms are prone to form segregations along the grain boundaries and triple junctions during SPD [20, 26, 27]. It should be noted that the impurities typically present in pure light metals also contribute to strengthening.

Precipitation strengthening. This type of strengthening is realized in metals and alloys containing second-phase precipitates that act as obstacles to dislocation glide. The strengthening effect is determined by the size and shape of the precipitates, their volume fraction, coherency of the matrix/precipitate interface, and spatial distribution of the precipitates [3].

Dislocation strengthening. The high density of dislocations in the grain/subgrain interior (statistically stored dislocations) and/or dislocation boundaries (geometrically necessary dislocations) can provide significant amount of strengthening in light metals [28]. Thus, this type of strengthening is typical for the SPD-processed light metals with a microstructure containing high dislocation density and a high volume fraction of LAGBs, which is usually formed after SPD at low temperatures to yield a low amount of strain (ϵvM ∼ 1…2) [21, 29].

Grain size hardening. The grain boundaries are major obstacles for the movement of dislocations; therefore, the distance traveled by a dislocation before reaching the grain boundary decreases with decreasing grain size, according to the well-known Hall–Petch law [30, 31]. There can be some amount of deviation from the Hall–Petch law if the grain size is reduced below 100 nm due to activation of the grain boundary–mediated mechanisms [32], presence of nonequilibrium grain boundaries introduced into material during SPD [33, 34], or effect of other strengthening mechanisms [13].

Crystallographic texture strengthening. SPD often results in the formation of a strong crystallographic texture in light metals, which can lead to significant anisotropy of the mechanical properties in Ti and Mg comprising HCP lattice [35–38]. The texture developed in these metals strongly depends on the SPD technique and processing parameters. For example, extrusion/drawing of pure Ti at low temperatures leads to the formation of α-fiber texture with the basal plane and c02-math-0001 being parallel to the extrusion/drawing direction [36, 39], while in pure Mg, the basal plane is inclined at 45 to the pressing direction after ECAP [37, 38].

It is well known that deformation twinning is suppressed in UFG Ti [40, 41] and Mg [42]. The main dislocation slip mode in pure Ti at room temperature is the prismatic c02-math-0002 slip system, which is followed, by the basal c02-math-0003 slip system, and pure Mg is deformed by basal slip c02-math-0004 . A low yield stress is typically observed if the main slip system is oriented at 45 with respect to the loading axis, resulting in the highest Schmid factor, S = 0.5 [43]. The yield stress increases with increasing/decreasing angle and can reach a maximum when the main slip system is parallel or perpendicular to the loading axis, which results in a very low Schmid factor [43]. Thus, the obtained texture can provide significant strengthening in both Ti and Mg.

A few strengthening mechanisms usually contribute to the strength of the NS light metals. The superposition law has been used to estimate the total contribution of several strengthening mechanisms, and the empirical equation has been expressed as

2.1 equation

where σ is the total yield strength, σss is the contribution of the solid solution hardening, σgs is the contribution of grain size hardening, σOr is the contribution of the Orowan strengthening, and the parameter n can vary between 1 and 2 [44]. It should be noted that Eq. (2.1) does not take into account the contribution of texture strengthening. However, in the strongly textured metals, grain size hardening is indirectly taken into account [45]

2.2 equation

In Eq. (2.2), the coefficient K can depend on the crystallographic texture of the metal and can be presented as

2.3 equation

where M is the Taylor factor, τCRSS is the critical resolved shear stress, and r is the distance from the nearest dislocation pile-up to the dislocation source in the adjacent grain [45].

As follows from Eq. (2.1), the total strength of the NS light metals should increase with increasing contribution of each strengthening mechanism. However, most often an increase in the contribution of one strengthening mechanism results in a decrease in the contribution of the other strengthening mechanism(s). For example, precipitation aging leads to increase in the Orowan strengthening, while solid solution hardening decreases due to purification of the matrix from solute atoms. Thus, the microstructural design for the generation of high-strength light metals is based on the concept of exploitation of the most efficient strengthening mechanisms at no or minimum compromise of the contribution of other mechanisms. It should also be noted that the interplay of different strengthening mechanisms has to be taken into account.

Table 2.2 shows the data on the mechanical properties of some of the NS light metals exhibiting superior strength. It was demonstrated that grain size hardening can be the most efficient strengthening mechanism in light metals [13, 21, 36, 51, 52], and a intelligent microstructural design is vital in order to gain the maximum contributions from the other strengthening mechanisms in order to achieve high mechanical strength. However, as already mentioned, , the texture effects can play very important role in the deformation behavior of the HCP metals. Figure 2.2a illustrates the effect of crystallographic texture on the mechanical behavior of UFG pure Ti with an average grain size of 200 nm. It is seen that the material exhibits extremely high tensile yield strength of 1250 MPa along the rod axis, whereas its yield strength in the transverse direction is just half of it although it is still higher as compared to its CG counterpart (Table 2.2, Figure 2.2a). A significant effect of orientation on the mechanical behavior is also clearly seen in Figure 2.2a. This significant anisotropy of the mechanical properties and mechanical behavior is rationalized on the basis of the strong α-fiber texture with the basal plane parallel to the rod axis (Figure 2.2b) [36].

Table 2.2 Data on the mechanical properties of some light metals and alloys exhibiting high mechanical strength

nfgz002

Figure 2.2 (a) Engineering stress–strain curves from tensile testing of UFG Ti (Grade 4) produced via complex SPD method (ECAP–swaging–drawing). (b) Pole figures from texture measurements of the material. LD – longitudinal direction; TD – transverse direction.

The images are reproduced from Ref. [36] with permission of the publisher.

Another good illustration of the texture effect on the mechanical properties and behavior of the AZ31B Mg alloy is presented in Figure 2.3a. This material was subjected to multitemperature ECAP with varying extrusion speed [23]. A significant texture-driven anisotropy was observed. The compression samples in the longitudinal (normal to the extrusion direction) and extrusion directions exhibited the lowest mechanical strength because the basal slip was eased in the grains, whereas the 37 ° samples demonstrated the highest